Cobalt based alloy additive manufactured article, cobalt based alloy product, and method for manufacturing same

ABSTRACT

There is provided an additive manufactured (AM) article formed of a Co based alloy having a composition comprising: in mass %, 0.08-0.25% C; 0.1% or less B; 10-30% Cr; 30% or less in total of Fe and Ni, the Fe being 5% or less; 5-12% in total of W and/or Mo; 0.5-2% in total of Ti, Zr, Nb and Ta; 0.5% or less Si; 0.5% or less Mn; 0.003-0.04% N; and the balance being Co and impurities. The AM article comprises crystal grains with an average size of 10-100 μm. In the crystal grains, segregation cells with an average size of 0.15-1.5 μm are formed, in which components constituting an MC type carbide phase comprising the Ti, Zr, Nb and/or Ta are segregated in boundary regions of the cells, and/or grains of the MC type carbide phase are precipitated at an average intergrain distance of 0.15-1.5 μm.

CLAIM OF PRIORITY

This application is a continuation of U.S. patent application Ser. No.16/124,978, filed Sep. 7, 2018, which is based upon and claims thebenefit of priority from Japanese patent application serial no.2017-172999 filed on Sep. 8, 2017, the content of which is herebyincorporated by reference into this application.

FIELD OF THE INVENTION

The present invention relates to cobalt based alloy articles havingexcellent mechanical properties and, in particular, to a cobalt basedalloy additive manufactured article, a cobalt based alloy product basedon the article, and a method for manufacturing the article and theproduct.

DESCRIPTION OF RELATED ART

Cobalt (Co) based alloy articles, along with nickel (Ni) based alloyarticles, are representative heat resistant alloy materials. Alsoreferred to as super alloys, they are widely used for high temperaturemembers of turbines (e.g. gas turbines, steam turbines, etc.). AlthoughCo based alloy articles are higher in material costs than Ni based alloyarticles, they have been used for applications such as turbine statorblades and combustor members because of their excellence in corrosionresistance and abrasion resistance, and their ease of solid solutionstrengthening.

In Ni based alloy materials, various improvements that have been made sofar in composition and manufacturing processes of heat resistant alloymaterials have led to the development of strengthening through γ′ phase(e.g. Ni₃(Al, Ti) phase) precipitation, which has now become mainstream.On the other hand, in Co based alloy materials, an intermetalliccompound phase that contributes to improving mechanical properties, likethe γ′ phase in Ni based alloy materials, hardly precipitates, which hasprompted research on carbide phase precipitation strengthening.

For example, JP Shou 61 (1986)-243143 A discloses a Co basedsuperplastic alloy made up of a Co based alloy matrix having a crystalgrain size of equal to or less than 10 μm and carbide grains in agranular form or a particulate form having a grain size of 0.5 to 10 μmprecipitated in the matrix. The Co based alloy includes 0.15 to 1 wt. %of C, 15 to 40 wt. % of Cr, 3 to 15 wt. % of W or Mo, 1 wt. % or less ofB, 0 to 20 wt. % of Ni, 0 to 1.0 wt. % of Nb, 0 to 1.0 wt. % of Zr, 0 to1.0 wt. % of Ta, 0 to 3 wt. % of Ti, 0 to 3 wt. % of Al, and the balanceof Co. According to JP Shou 61 (1986)-243143 A, there can be provided aCo based superplastic alloy that exhibits superplasticity accompanyingwith an elongation of equal to or more than 70% even in a lowtemperature range (e.g. at 950° C.), and is capable of being formed intoan article with a complicated shape by plastic working such as forging.

JP Hei 7 (1995)-179967 A discloses a Co based alloy that is excellent incorrosion resistance, abrasion resistance, and high temperaturestrength. The alloy includes 21 to 29 wt. % of Cr, 15 to 24 wt. % of Mo,0.5 to 2 wt. % of B, 0.1 or more and less than 0.5 wt. % of Si, morethan 1 and equal to or less than 2 wt. % of C, 2 wt. % or less of Fe, 2wt. % or less of Ni, and the balance of substantially Co. According toJP Hei 7 (1995)-179967 A, the Co based alloy has a composite structurein which a molybdenum boride and a chromium carbide are relativelyfinely dispersed in a quaternary alloy phase of Co, Cr, Mo and Si andexhibits excellent corrosion resistance, abrasion resistance, and highstrength.

Meanwhile, in recent years, three dimensional shaping technology (theso-called 3D printing) such as additive manufacturing or AM has receivedmuch attention as a technique for manufacturing finished products with acomplicated shape by near net shaping. To apply the three dimensionalshaping technology to heat resistant alloy components, vigorous researchand development activities are currently being carried out.

For example, JP 2016-535169 A (WO 2015/073081 A1) discloses a method ofproducing layers including the steps of: (a) providing a source of agranular composite powder having a porosity of 20% or less; (b)depositing a first portion of said powder onto a target surface; (c)depositing energy into the powder of said first portion under conditionsthat said energy causes sintering, fusing or melting of the first powderportion so as to create a first layer; (d) depositing a second portionof powder onto said first layer; and (e) depositing energy into thepowder of said second portion under conditions that said energy causessintering, fusing or melting of the second powder portion so as tocreate a second layer. In the method, the energy is supplied by a laserpositioned on said first layer.

JP 2016-535169 A (WO 2015/073081 A1) teaches as follows: Selective lasermelting (SLM) or direct metal laser melting (DMLM) uses laser to make amaterial a full melt. Full melting is typically useful for amonomaterial (e.g. pure titanium or a single alloy such as Ti-6Al-4V),as there is just one melting point. By contrast, selective lasersintering (SLS) and direct metal laser sintering (DMLS) are essentiallythe same thing, and SLS/DMLS is used to apply processes to a variety ofmaterials—multiple metals, alloys, or combinations of alloys and othermaterials such as plastics, ceramics, polymers, carbides or glasses.Meanwhile, sintering is apart from melting, and a sintering process doesnot fully melt a material but heats it to the point that the materialcan fuse together on a molecular level.

Since the 3D printing is capable of directly forming even components ofcomplicated shape, manufacturing of turbine high temperature componentsby the 3D printing is very attractive in terms of reduction ofmanufacturing work time and improvement of manufacturing yield (i.e.reduction of manufacturing cost).

Co based alloy materials such as the ones disclosed in JP Show 61(1986)-243143 A and JP Hei 7 (1995)-179967 A are thought to havemechanical properties superior to those of previous Co based alloymaterials. Unfortunately, however, their mechanical properties areinferior to those of precipitation-strengthened Ni based alloy materialsof recent years. Therefore, many studies on additive manufacturingarticles (AM articles) for use as turbine high temperature componentsare currently directed toward precipitation-strengthened Ni based alloymaterials.

However, AM articles of the precipitation-strengthened Ni based alloysare prone to have problems such as generation of the γ′ phase, which isthe core of their mechanical properties, being hindered and internaldefects occurring in the finished products. As a result, expectedmechanical properties have not been sufficiently achieved. This isattributable to the fact that current precipitation-strengthened Nibased alloy materials used for turbine high temperature components havebeen optimized through melting and forging processes under high vacuum,and therefore oxidation and nitriding of the Al component and the Ticomponent, which constitute the γ′ phase, easily occur at the stages ofpreparing alloy powder for AM and performing AM.

On the other hand, manufacturing the Co based alloy materials such asthe ones disclosed in JP show 61 (1986)-243143 A and JP Hei 7(1995)-179967 A does not require precipitation of an intermetalliccompound phase such as the γ′ phase as in Ni based alloy materials, soCo based alloy materials do not contain plenty of Al or Ti, which iseasily oxidized. This means melting and forging processes in the airatmosphere are available for their manufacturing. Therefore, such Cobased alloy materials are considered to be advantageous in manufacturingof alloy powder for AM and manufacturing of AM articles. Also, the Cobased alloy materials have advantages with corrosion resistance andabrasion resistance comparable to or superior to those of Ni based alloymaterials.

However, as mentioned above, conventional Co based alloy materials havedisadvantages of mechanical properties inferior to those of γ′ phaseprecipitation-strengthened Ni based alloy materials. In other words, ifa Co based alloy material could achieve mechanical properties comparableto or superior to those of γ′ phase precipitation-strengthened Ni basedalloy materials (e.g. a creep temperature endurable for 100,000 hours ata stress of 58 MPa being equal to or higher than 875° C., a tensilestrength being equal to or more than 500 MPa at room temperature, etc.),AM articles of the Co based alloy material would become highlyattractive turbine high temperature components.

SUMMARY OF THE INVENTION

The present invention was made in view of the foregoing and has anobjective to provide an additive manufactured article of a Co basedalloy material having mechanical properties comparable to or superior tothose of precipitation-strengthened Ni based alloy materials, and a Cobased alloy product based on the additive manufacturing article. Also,another objective is to provide a method for manufacturing the Co basedalloy additive manufactured article and the Co based alloy product.

(I) According to one aspect of the present invention, there is providedan additive manufactured article formed of a cobalt based alloy. Thecobalt based alloy has a chemical composition including: 0.08 to 0.25mass % of carbon (C); 0.1 mass % or less of boron (B); 10 to 30 mass %of chromium (Cr); 30 mass % or less in total of iron (Fe) and nickel(Ni), the Fe being in an amount of 5 mass % or less; 5 to 12 mass % intotal of tungsten (W) and/or molybdenum (Mo); 0.5 to 2 mass % in totalof titanium (Ti), zirconium (Zr), niobium (Nb) and tantalum (Ta); 0.5mass % or less of silicon (Si); 0.5 mass % or less of manganese (Mn);0.003 to 0.04 mass % of nitrogen (N); and the balance being cobalt (Co)and impurities. The additive manufactured article is a polycrystallinebody comprising crystal grains with an average crystal grain size of 10to 100 μm. In the crystal grains of the polycrystalline body,segregation cells with an average size of 0.15 to 1.5 μm are formed, inwhich components constituting an MC type carbide phase comprising theTi, Zr, Nb and/or Ta are segregated in boundary regions of thesegregation cells, and/or grains of the MC type carbide phase areprecipitated at an average intergrain distance of 0.15 to 1.5 μm.

In the above Co based alloy additive manufactured article (I) of theinvention, the following changes and modifications can be made.

(i) The chemical composition of the Co based alloy may include 0.01 to 1mass % of the Ti, 0.05 to 1.5 mass % of the Zr, 0.02 to 1 mass % of theNb, and 0.05 to 1.5 mass % of the Ta; and

(ii) The chemical composition of the Co based alloy may include, as theimpurities, 0.5 mass % or less of aluminum (Al) and 0.04 mass % or lessof oxygen (O).

(II) According to another aspect of the invention, there is provided aproduct formed of a Co based alloy, the Co based alloy having a chemicalcomposition including: 0.08 to 0.25 mass % of C; 0.1 mass % or less ofB; 10 to 30 mass % of Cr; 30 mass % or less in total of Fe and Ni, theFe being in an amount of 5 mass % or less; 5 to 12 mass % in total of Wand/or Mo; 0.5 to 2 mass % in total of Ti, Zr, Nb and Ta; 0.5 mass % orless of Si; 0.5 mass % or less of Mn; 0.003 to 0.04 mass % of N; and thebalance being Co and impurities. The product is a polycrystalline bodycomprising crystal grains with an average crystal grain size of 20 to145 μm. In the crystal grains of the polycrystalline body of product,grains of an MC type carbide phase comprising the Ti, Zr, Nb and/or Taare precipitated at an average intergrain distance of 0.15 to 1.5 μm.

In the above Co based alloy product (II) of the invention, the followingchanges and modifications can be made.

(iii) The product may exhibit a 0.2% proof stress of 500 MPa or more atroom temperature and a creep temperature endurable for 100,000 hours ata stress of 58 MPa being 875° C. or higher;

(iv) The chemical composition of the Co based alloy may include 0.01 to1 mass % of the Ti, 0.05 to 1.5 mass % of the Zr, 0.02 to 1 mass % ofthe Nb, and 0.05 to 1.5 mass % of the Ta;

(v) The chemical composition of the Co based alloy may include, as theimpurities, 0.5 mass % or less of Al and 0.04 mass % or less of O;

(vi) The product may be a turbine high temperature component; and

(vii) The turbine high temperature component may be a turbine statorblade or a combustor nozzle.

(III) According to still another aspect of the invention, there isprovided a method for manufacturing the above Co based alloy additivemanufactured article. The manufacturing method includes: an alloy powderpreparation step of preparing a Co based alloy powder having thechemical composition, the Co based alloy powder having a predeterminedparticle size; and a selective laser melting step of forming an additivemanufactured article, the step comprising alternate repetition of analloy powder bed preparation substep of laying the Co based alloy powdersuch that it forms an alloy powder bed having a predetermined thicknessand a laser melting solidification substep of irradiating apredetermined region of the alloy powder bed with a laser beam tolocally melt and rapidly solidify the Co based alloy powder in theregion, the laser beam having a predetermined output power and apredetermined scanning speed. The predetermined thickness of the alloypowder bed h (unit: μm), the predetermined output power of the laserbeam P (unit: W), and the predetermined scanning speed of the laser beamS (unit: mm/s) are controlled to satisfy the following formulas:“15<h<150” and “67×(P/S)−3.5<h<2222×(P/S)+13”.

In the above manufacturing method of the Co based alloy additivemanufactured article (III) of the invention, the following changes andmodifications can be made.

(viii) The alloy powder preparation step may include an alloy powderclassification substep of regulating the particle size of the Co basedalloy powder to 5 to 100 μm.

(IV) According to still another aspect of the invention, there isprovided a method for manufacturing the above Co based alloy product.The manufacturing method includes: an alloy powder preparation step ofpreparing a cobalt based alloy powder having the chemical composition,the cobalt based alloy powder having a predetermined particle size; aselective laser melting step of forming an additive manufacturedarticle, the step comprising alternate repetition of an alloy powder bedpreparation substep of laying the Co based alloy powder such that itforms an alloy powder bed having a predetermined thickness and a lasermelting solidification substep of irradiating a predetermined region ofthe alloy powder bed with a laser beam to locally melt and rapidlysolidify the Co based alloy powder in the region, the laser beam havinga predetermined output power and a predetermined scanning speed; and asolution heat treatment step of subjecting the additive manufacturedarticle to a solution treatment at temperatures ranging from 1,100 to1,200° C. with a holding duration of 0.5 to 10 hours. The predeterminedthickness of the alloy powder bed h (unit: μm), the predetermined outputpower of the laser beam P (unit: W), and the predetermined scanningspeed of the laser beam S (unit: mm/s) are controlled to satisfy thefollowing formulas: “15<h<150” and “67×(P/S)−3.5<h<2222×(P/S)+13”.

In the above manufacturing method of the Co based alloy product (IV) ofthe invention, the following changes and modifications can be made.

(ix) The alloy powder preparation step may include an alloy powderclassification substep of regulating the particle size of the cobaltbased alloy powder to 5 to 100 μm; and

(x) The manufacturing method may further include an aging heat treatmentstep of subjecting the additive manufactured article to an agingtreatment at temperatures ranging from 750 to 1,000° C. with a holdingduration of 0.5 to 10 hours after the solution heat treatment step.

Advantages of the Invention

According to the present invention, there can be provided an additivemanufactured article of a Co based alloy material having mechanicalproperties comparable to or superior to those ofprecipitation-strengthened Ni based alloy materials, a Co based alloyproduct based on the additive manufacturing article, and a method formanufacturing the Co based alloy additive manufactured article and theCo based alloy product.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a flow diagram showing an exemplary process of a method formanufacturing a Co based alloy product according to an embodiment of thepresent invention;

FIG. 2 is an example of a scanning electron microscope (SEM) image of aCo based alloy AM article according to an embodiment of the invention;

FIG. 3 is a schematic illustration of a perspective view showing aturbine stator blade which is a Co based alloy product as a turbine hightemperature component according to an embodiment of the invention;

FIG. 4 is a schematic illustration of a cross-sectional view showing agas turbine equipped with a Co based alloy product according to anembodiment of the invention;

FIG. 5 is an SEM image showing an exemplary microstructure of a Co basedalloy AM article formed by LMD;

FIG. 6 is an SEM image showing an exemplary microstructure of a Co basedalloy cast article formed by precision casting;

FIG. 7 is a graph showing an exemplary relationship between an averagesize of segregation cells of Co based alloy AM articles and a 0.2% proofstress of Co based alloy products; and

FIG. 8 shows exemplary SLM conditions to obtain a Co based alloy AMarticle according to an embodiment of the invention, indicating arelationship between a thickness of an alloy powder bed and a local heatinput.

DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

[Basic Concept of the Present Invention]

As mentioned before, various research and development activities havebeen carried out on strengthening of Co based alloy materials throughcarbide phase precipitation. Carbide phases that contribute toprecipitation strengthening include MC type carbide phases of Ti, Zr, Nband Ta, and complex carbide phases of these metallic elements.

The C component, which is indispensable in formation of a carbide phasewith each of Ti, Zr, Nb and Ta, tends to segregate significantly atfinal solidification portions (e.g. dendrite boundaries, crystal grainboundaries, etc.) at the melting and solidification stages of the Cobased alloy. So, in conventional Co based alloy materials, the carbidephase grains precipitate along the dendrite boundaries and crystal grainboundaries in the matrix. In a general cast material of a Co basedalloy, for example, the average spacing between dendrite boundaries andthe average crystal grain size are on the order of 10¹ to 10² μm, andtherefore the average spacing between carbide phase grains is also onthe order of 10¹ to 10² μm. Furthermore, even with the relatively fastsolidification rate of laser welding, for example, the average spacingbetween carbide phase grains at the solidified portions is around 5 μm.

Precipitation strengthening in alloys is generally known to be inverselyproportional to the average spacing between precipitates, and it isconsidered that precipitation strengthening is effective only when theaverage spacing between precipitates is around 2 μm or less. However,with the above-mentioned conventional technology, the average spacingbetween precipitates has not reached this level in a Co based alloymaterial, and sufficient precipitation strengthening effect has not beenachieved. In other words, with the conventional technology, it has beendifficult to finely and dispersedly precipitate carbide phase grainsthat might contribute to strengthening alloys. This would be the mainfactor behind the fact that Co based alloy materials have been said tohave mechanical properties inferior to those ofprecipitation-strengthened Ni based alloy materials.

Meanwhile, another carbide phase that can precipitate in Co based alloysis the Cr carbide phase. Since the Cr component is highly solid solubleto the Co based alloy matrix and hardly segregate, the Cr carbide phasecan be dispersedly precipitated in the matrix crystal grains. However,the Cr carbide phase is poor in lattice matching with the Co based alloymatrix crystals and is known to be not so effective as aprecipitation-strengthening phase.

The present inventors thought that if they were able to dispersedlyprecipitate carbide phase grains contributing to precipitationstrengthening in the matrix crystal grains, they would be able todramatically improve mechanical properties of Co based alloy materials.Considering the inherent excellent corrosion resistance and abrasionresistance of Co based alloy materials, they would be able to provide aheat-resistant alloy material that would surpassprecipitation-strengthened Ni based alloy materials.

In order to obtain such a Co based alloy material, the inventorsconducted intensive research on alloy compositions and manufacturingmethods. As a result, they have found that it is possible to dispersedlyprecipitate carbide phase grains contributing to alloy strengthening inthe matrix crystal grains of a Co based alloy material by optimizing thealloy composition and controlling the amount of heat input for localmelting and rapid solidification in a manufacturing method including AM(in particular, selective laser melting). The present invention was madebased on this finding.

Preferred embodiments of the present invention will be hereinafterdescribed with reference to the accompanying drawings. However, theinvention is not limited to the specific embodiments described below,and various combinations with known art and modifications based on knownart are possible without departing from the spirit and the scope of theinvention.

[Method for Manufacturing Co Based Alloy Additive ManufacturingArticle/Co Based Alloy Product]

FIG. 1 is a flow diagram showing an exemplary process of a method formanufacturing a Co based alloy product according to an embodiment of theinvention. As shown in FIG. 1, the method for manufacturing a Co basedalloy product roughly includes: an alloy powder preparation step S1 ofpreparing a Co based alloy powder; a selective laser melting step S2 offorming the prepared Co based alloy powder into an AM article with adesired shape; a solution heat treatment step S3 of subjecting theformed AM article to a solution treatment; and an aging heat treatmentstep S4 of subjecting the solution heat-treated AM article to an agingtreatment. The AM article obtained by the selective laser melting stepS2 is a Co based alloy additive manufactured article according to anembodiment of the invention.

Each step will be hereinafter described in more detail.

(Alloy Powder Preparation Step)

In the step S1, a Co based alloy powder having a predetermined chemicalcomposition is prepared. The chemical composition preferably includes:0.08 to 0.25 mass % of C; 0.1 mass % or less of B; 10 to 30 mass % ofCr; 30 mass % or less in total of Fe and Ni with Fe being in an amountof 5 mass % or less; 5 to 12 mass % in total of W and/or Mo; 0.5 to 2mass % in total of Ti, Zr, Nb and Ta; 0.5 mass % or less of Si; 0.5 mass% or less of Mn; 0.003 to 0.04 mass % of N; and the balance being Co andimpurities.

C: 0.08 to 0.25 Mass %

The Co component is an important component that constitutes an MC typecarbide phase to serve as a precipitation strengthening phase (this MCtype carbide phase may be referred to as “carbide phase of Ti, Zr, Nband/or Ta or “reinforcing carbide phase”). The content of the Ccomponent is preferably 0.08 to 0.25 mass %, more preferably 0.1 to 0.2mass %, and even more preferably 0.12 to 0.18 mass %. When the C contentis less than 0.08 mass %, the amount of precipitation of the reinforcingcarbide phase is insufficient, resulting in an insufficient effect ofimproving the mechanical properties. By contrast, when the C content isover 0.25 mass %, the alloy material becomes excessively hard, whichleads to deteriorated ductility and toughness.

B: 0.1 Mass % or Less

The B component contributes to improving bondability between crystalgrain boundaries (the so-called grain boundary strengthening). Althoughthe B is not an essential component, when it is contained in the alloy,the content of the B component is preferably 0.1 mass % or less and morepreferably 0.005 to 0.05 mass %. When the B component is over 0.1 mass%, cracking (e.g. solidification cracking) is prone to occur duringformation of the AM article.

Cr: 10 to 30 Mass %

The Cr component contributes to improving corrosion resistance andoxidation resistance. The content of the Cr component is preferably 10to 30 mass % and more preferably 10 to 25 mass %. In the case where acorrosion resistant coating layer is provided on the outermost surfaceof the Co based alloy product, the content of the Cr component is evenmore preferably 10 to 18 mass %. When the Cr content is less than 10mass %, the corrosion resistance and the oxidation resistance areinsufficient. Meanwhile, when the Cr content is over 30 mass %, thebrittle σ phase and/or the excessive amount of Cr carbide phase aregenerated, resulting in deteriorated mechanical properties (i.e.toughness, ductility, strength, etc.).

Ni: 30 Mass % or Less

Being similar to Co component in properties but less expensive than Co,the Ni component may be used to replace part of the Co component.Although the Ni is not an essential component, when it is contained inthe alloy, the content of the Ni component is preferably 30 mass % orless, more preferably 20 mass % or less, and more preferably 5 to 15mass %. When the Ni content is over 30 mass %, the abrasion resistanceand the local stress resistance, which are characteristics of Co basedalloys, deteriorate. This is attributable to the difference in stackingfault energy between Co and Ni.

Fe: 5 Mass % or Less

Being much less expensive than Ni and similar to Ni component inproperties, the Fe component may be used to replace part of the Nicomponent. The total content of the Fe and Ni is preferably 30 mass % orless, more preferably 20 mass % or less, and even more preferably 5 to15 mass %. Although the Fe is not an essential component, when it iscontained in the alloy, the content of the Fe component is preferably 5mass % or less and more preferably 3 mass % or less. When the Fe contentis over 5 mass %, the corrosion resistance and mechanical propertiesdeteriorate.

W and/or Mo: 5 to 12 Mass % in Total

The W component and the Mo component contribute tosolution-strengthening the matrix. The total content of the W componentand/or the Mo component is preferably 5 to 12 mass % and more preferably7 to 10 mass %. When the total content of the W component and the Mocomponent is less than 5 mass %, the solution strengthening of thematrix is insufficient. In contrast, when the total content of the Wcomponent and the Mo component is over 12 mass %, the brittle σ phasetends to be generated easily, resulting in deteriorated mechanicalproperties (i.e. toughness, ductility, etc.).

Re: 2 Mass % or Less

The Re component contributes to solution-strengthening the matrix andimproving corrosion resistance. Although the Re is not an essentialcomponent, when it is contained in the alloy to replace part of the Wcomponent or the Mo component, the content of the Re component ispreferably 2 mass % or less and more preferably 0.5 to 1.5 mass %. Whenthe Re content is over 2 mass %, the functions and effects of the Recomponent become saturated, and the material costs become too high.

Ti, Zr, Nb, and Ta: 0.5 to 2 Mass % in Total

The Ti component, the Zr component, the Nb component, and the Tacomponent are important components that constitute the reinforcingcarbide phase (MC type carbide phase). The total content of the Ti, Zr,Nb and Ta components is preferably 0.5 to 2 mass % and more preferably0.5 to 1.8 mass %. When the total content is less than 0.5 mass %, theamount of precipitation of the reinforcing carbide phase isinsufficient, and, as a result, the effect of improving the mechanicalproperties is insufficient. By contrast, when the total content is over2 mass %, the mechanical properties deteriorate due to coarsening of thegrains of the reinforcing carbide phase, accelerated generation of abrittle phase (e.g. σ phase), generation of grains of an oxide phasethat does not contribute to precipitation strengthening, etc.

More specifically, the Ti content is preferably 0.01 to 1 mass % andmore preferably 0.05 to 0.8 mass %. The Zr content is preferably 0.05 to1.5 mass % and more preferably 0.1 to 1.2 mass %. The Nb content ispreferably 0.02 to 1 mass % and more preferably 0.05 to 0.8 mass %. TheTa content is preferably 0.05 to 1.5 mass % and more preferably 0.1 to1.2 mass %.

Si: 0.5 Mass % or Less

The Si component serves as a deoxidant agent and contributes toimproving the mechanical properties. Although the Si is not an essentialcomponent, when it is contained in the alloy, the content of the Sicomponent is preferably 0.5 mass % or less and more preferably 0.01 to0.3 mass %. When the Si content is over 0.5 mass %, coarse grains of anoxide (e.g. SiO₂) are generated, which causes deterioration of themechanical properties.

Mn: 0.5 Mass % or Less

The Mn component serves as a deoxidant agent and a desulfurizing agentand contributes to improving the mechanical properties and the corrosionresistance. Although the Mn is not an essential component, when it iscontained in the alloy, the content of the Mn component is preferably0.5 mass % or less and more preferably 0.01 to 0.3 mass %. When the Mncontent is over 0.5 mass %, coarse grains of a sulfide (e.g. MnS) aregenerated, which causes deterioration of the mechanical properties andthe corrosion resistance.

N: 0.003 to 0.04 Mass %

The N component contributes to stabilizing the generation of thereinforcing carbide phase. The content of the N component is preferably0.003 to 0.04 mass %, more preferably 0.005 to 0.03 mass %, and evenmore preferably 0.007 to 0.025 mass %. When the N content is less than0.003 mass %, the functions and effects of the N component areinsufficient. Meanwhile, when the N content is over 0.04 mass %, coarsegrains of a nitride (e.g. Cr nitride) are generated, which causesdeterioration of the mechanical properties.

Balance: Co Component and Impurities

The Co component is one of the key components of the alloy and itscontent is the largest of all the components. As mentioned above, Cobased alloy materials have the advantages of having corrosion resistanceand abrasion resistance comparable to or superior to those of Ni basedalloy materials.

The Al component is one of the impurities of the alloy and is not to beintentionally included in the alloy. However, an Al content of 0.3 mass% or less is acceptable as it does not have any serious negativeinfluence on the mechanical properties of the Co based alloy product.When the Al content is over 0.3 mass %, coarse grains of an oxide ornitride (e.g. Al₂O₃ or AlN) are generated, which causes deterioration ofthe mechanical properties.

The O component is also one of the impurities of the alloy and is not tobe intentionally included in the alloy. However, an O content of 0.04mass % or less is acceptable as it does not have any serious negativeinfluence on the mechanical properties of the Co based alloy product.When the O content is over 0.04 mass %, coarse grains of each oxide(e.g. Ti oxide, Zr oxide, Al oxide, Fe oxide, Si oxide, etc.) aregenerated, which causes deterioration of the mechanical properties.

In this step S1, there is no particular limitation on the method andtechniques for preparing the Co based alloy powder, and any conventionalmethod and technique may be used. For example, a master ingotmanufacturing substep S1 a of manufacturing a master ingot by mixing,melting, and casting the raw materials such that the ingot has a desiredchemical composition and an atomization substep S1 b of forming thealloy powder from the master ingot may be performed. Also, there is noparticular limitation on the atomization method, and any conventionalmethod and technique may be used. For example, gas atomizing orcentrifugal force atomizing, by which spherical particles of high purityand homogeneous composition can be obtained, may be preferably used.

For ease of handling and ease of filling the alloy powder bed in thefollowing selective laser melting step S2, the particle size of thealloy powder is preferably 5 to 100 μm, more preferably 10 to 70 μm, andeven more preferably 10 to 50 μm. When the particle size of the alloypowder is less than 5 μm, fluidity of the alloy powder decreases in thefollowing step S2 (i.e. formability of the alloy powder bed decreases),which causes deterioration of shape accuracy of the AM article. Incontrast, when the particle size of the alloy powder is over 100 μm,controlling the local melting and rapid solidification of the alloypowder bed in the following step S2 becomes difficult, which leads toinsufficient melting of the alloy powder and increased surface roughnessof the AM article.

In view of the above, an alloy powder classification substep S1 c ispreferably performed so as to regulate the alloy powder particle size to5 to 100 μm. In the present invention, when the particle sizedistribution of the alloy powder manufactured in the atomization substepS1 b is observed to fall within the desired range, it is assumed thatthe substep S1 c has been performed.

(Selective Laser Melting Step)

In the step S2, the prepared Co based alloy powder is formed into an AMarticle having a desired shape by selective laser melting (SLM). Morespecifically, this step comprises alternate repetition of an alloypowder bed preparation substep S2 a and a laser melting solidificationsubstep S2 b. In the step S2 a, the Co based alloy powder is laid suchthat it forms an alloy powder bed having a predetermined thickness, andin the step S2 b, a predetermined region of the alloy powder bed isirradiated with a laser beam to locally melt and rapidly solidify the Cobased alloy powder in the region.

In this step S2, in order to obtain a finished Co based alloy producthaving a desired microstructure (a microstructure in which thereinforcing carbide phase grains are dispersedly precipitated in thematrix crystal grains), the microstructure of the AM article, which is aprecursor of the finished product, is controlled by controlling thelocal melting and the rapid solidification of the alloy powder bed.

More specifically, the thickness of the alloy powder bed h (unit: μm),the output power of the laser beam P (unit: W), and the scanning speedof the laser beam S (unit: mm/s) are preferably controlled to satisfythe following formulas: “15<h<150” and “67×(P/S)−3.5<h<2222×(P/S)+13”.When these formulas are not satisfied, an AM article having a desiredmicrostructure cannot be obtained. This step S2 makes it possible toobtain a Co based alloy AM article according to an embodiment of theinvention.

While the output power P and the scanning speed S of the laser beambasically depend on configurations of the laser apparatus, they may bedetermined so as to satisfy the following formulas: “10≤P≤1000” and“10≤S≤7000”.

(Co Based Alloy Additive Manufactured Article)

FIG. 2 is an example of a scanning electron microscope (SEM) image of aCo based alloy AM article according to an embodiment of the invention.As shown in FIG. 2, the Co based alloy AM article according to theembodiment of the invention has a unique microstructure that has neverbeen seen before.

The AM article is a polycrystalline body with an average crystal grainsize of 10 to 100 μm. In the crystal grains of the polycrystalline body,segregation cells (also capable of being referred to segregationmicrocells) with an average size of 0.15 to 1.5 μm are formed. Also,grains of the reinforcing carbide phase are precipitated at an averagespacing of 0.15 to 1.5 μm. In the present invention, the size ofsegregation cells is defined as the average of the long diameter and theshort diameter.

A more detailed microstructure observation by transmission electronmicroscopy-energy dispersive X-ray spectrometry (TEM-EDX) has revealedthat the components constituting the reinforcing carbide phase (Ti, Zr,Nb, Ta, and C) segregate in the boundary regions between the neighboringsegregation cells (i.e. in outer peripheral regions of microcells,similar to cell walls). It has also been observed that grains of thereinforcing carbide phase are precipitated at part of the triplepoints/quadruple points in the boundary regions among these segregationcells.

(Solution Heat Treatment Step)

In the step S3, the formed Co based alloy AM article is subjected to asolution treatment. The heat treatment is preferably performed attemperatures ranging from 1,100 to 1,200° C. with a holding duration of0.5 to 10 hours. There is no particular limitation on a cooling methodafter the heat treatment, and oil cooling, water cooling, air cooling,or furnace cooling may be used.

This solution treatment allows the crystal grains of the matrix of theAM article to recrystallize, thereby reducing the internal strain in thematrix crystal grains of the AM article that has occurred during rapidcooling solidification. It is preferable that the average crystal grainsize of the matrix crystal grains be controlled to 20 to 145 μm. Whenthe average crystal grain size is less than 20 μm or over 145 μm, thefinished Co based alloy product does not exhibit sufficient creepproperties.

In addition, interestingly enough, it has been found that as the matrixcrystal grains recrystallize, the components segregated in the boundaryregions of the segregation cells start to aggregate to form thereinforcing carbide phase, and as a result, the segregation cellsdisappear (more specifically, they become unobservable by scanningelectron microscopy, or SEM). The aggregation points where they begin toform the reinforcing carbide phase are thought to be at triple pointsand quadruple points of the former segregation cell boundaries, whichcauses the fine dispersion of the reinforcing carbide phase throughoutthe matrix crystal grains (within each crystal grain and on the crystalgrain boundaries).

By successfully controlling the temperature and the holding duration ofthe solution heat treatment, the reinforcing carbide phase beginning toform can be grown into grain form without being aggregated/coarsenedexcessively. In such a case, the Co based alloy product obtained throughthis step S3 may be considered to be a finished product. However, interms of preventing from the excessive coarsening of the matrix crystalgrains (in other words, in terms of manufacturing stability and yield),it is more preferable that the following aging heat treatment step S4 beperformed.

(Aging Heat Treatment Step)

In the step S4, the solution-treated Co alloy AM articles is subjectedto an aging treatment. The aging treatment is preferably performed attemperatures ranging from 750 to 1000° C. with a holding duration of 0.5to 10 hours. There is no particular limitation on the cooling methodafter the heat treatment, and oil cooling, water cooling, air cooling,or furnace cooling may be used.

The aging treatment allows the reinforcing carbide phase beginning toform in the solution heat treatment step S3 to grow into grain formwhile controlling excessive coarsening of the matrix crystal grains. TheCo based alloy product thus obtained has an average matrix crystal grainsize of 20 to 145 μm and includes grains of the reinforcing carbidephase finely and dispersedly precipitated in each crystal grain at anaverage intergrain distance of 0.15 to 1.5 μm. Naturally enough, grainsof the reinforcing carbide phase are dispersedly precipitated also onthe matrix crystal grain boundaries in the Co based alloy productaccording to the embodiment of the invention.

(Finishing Step) Where appropriate, other steps such as a step offorming a corrosion resistant coating layer and a surface finishingstep, not shown in FIG. 1, may be further performed on the Co basedalloy product obtained through the solution heat treatment step S3 orthe aging heat treatment step S4.

[Co Based Alloy Product]

FIG. 3 is a schematic illustration of a perspective view showing aturbine stator blade which is a Co based alloy product as a turbine hightemperature component according to an embodiment of the invention. Asshown in FIG. 3, the turbine stator blade 100 includes an inner ringside end wall 101, a blade part 102, and an outer ring side end wall103. Inside the blade part 102 is often formed a cooling structure. Asseen from FIG. 3, since the turbine stator blade 100 has a verycomplicated shape and structure, the technical significance of AMarticles manufactured by near net shaping and alloy products based onsuch AM articles is profound.

Meanwhile, in the case of a gas turbine for power generation with anoutput of around 30 MW, the length of the blade part 102 of the turbinestator blade 100 (i.e. distance between the end walls 101 and 103) isapproximately 170 mm.

FIG. 4 is a schematic illustration of a cross-sectional view showing agas turbine equipped with a Co based alloy product according to anembodiment of the invention. As shown in FIG. 4, the gas turbine 200roughly includes a compression part 210 for compressing intake air and aturbine part 220 for blowing combustion gas of a fuel on turbine bladesto obtain rotation power. The turbine high temperature componentaccording to the embodiment of the invention can be preferably used as aturbine nozzle 221 or the turbine stator blade 100 inside the turbinepart 220. Naturally enough, the turbine high temperature componentaccording to the embodiment of the invention is not limited to gasturbine applications but may be used for other turbine applications(e.g. steam turbines).

EXAMPLES

The present invention will be hereinafter described in more detail withexamples and comparative examples. It should be noted that the inventionis not limited to these examples.

[Experimental 1]

(Preparation of Alloy Powders IA-1 to IA-5 of Inventive Examples andAlloy Powders CA-1 to CA-5 of Comparative Examples)

Co based alloy powders having the chemical compositions shown in Table 1were prepared (the alloy powder preparation Step S1). More specifically,first, the master ingot manufacturing substep S1 a was performed, inwhich the raw materials were mixed and subjected to melting and castingby a vacuum high frequency induction melting method so as to form amater ingot (weight: approximately 2 kg) for each powder. Next, theatomization substep S1 b was performed to form each alloy powder. In thesubstep S1 b, each master ingot was remelted and subjected to gasatomizing in an argon gas atmosphere.

Then, each alloy powder thus obtained was subjected to the alloy powderclassification substep S1 c to control the particle size. At this point,each alloy powder was classified into an alloy powder with a particlesize of 10 to 25 μm (particle size S) and another alloy powder with aparticle size of 100 to 150 μm (particle size L).

TABLE 1 Chemical Compositions of Alloy Powders IA-1 to IA-5 of InventiveExamples and Alloy Powders CA-1 to CA-5 of Comparative Examples. AlloyChemical Composition (Mass %) Ti + Zr + Number C B Cr Ni Fe W Ti Zr NbTa Si Mn N Co Al O Nb + Ta IA-1 0.08 0.009 24.5  9.5 0.01 7.3 0.15 0.400.05 0.20 0.01 0.01 0.005 Bal. 0.01 0.005 0.80 IA-2 0.16 0.011 25.5 10.50.90 7.7 0.30 0.60 0.15 0.40 0.30 0.20 0.025 Bal. 0.05 0.020 1.45 IA-30.25 0.009 30.0 — — 5.0 0.01 0.30 0.05 0.10 0.05 0.01 0.005 Bal. — 0.0050.46 IA-4 0.16 0.010 25.0 10.0 0.02 7.5 0.25 0.05 0.09 0.30 0.01 0.020.010 Bal. — 0.010 0.69 IA-5 0.10 0.011 10.0 29.1 0.90 12.0 0.60 0.600.15 0.50 0.30 0.20 0.030 Bal. 0.05 0.030 1.85 CA-1 0.35 0.009 32.0  9.50.01 7.3 0.15 0.40 0.05 0.50 0.01 0.01 0.005 Bal. 0.01 0.005 1.10 CA-20.35 0.009 30.0 40.0 0.01 7.3 0.90 0.40 1.0  1.0  0.01 0.01 0.005 Bal.2.20 0.005 3.30 CA-3 0.40 0.010 29.0 10.0 0.20 7.5 0.20 0.10 0.10 — 0.100.02 0.001 Bal. — 0.015 0.40 CA-4 0.25 0.010 29.0 10.0 0.10 7.0 — — — —— 0.01 0.010 Bal. — 0.010 0 CA-5 0.11 0.002 22.0 23.0 0.01 14.0 0.010.01 — — 0.50 0.003 0.006 Bal. 0.01 0.008 0.02 “—” indicates that theelement was not intentionally included or not detected. “Bal.” indicatesinclusion of impurities other than Al and O.

As shown in Table 1, the alloy powders IA-1 to IA-5 have chemicalcompositions that satisfy the specifications of the invention. Incontrast, the alloy powder CA-1 has a C content and a Cr content thatfail to satisfy the specifications of the invention. The alloy powderCA-2 has a C content, an Ni content, and a total content of“Ti+Zr+Nb+Ta” that are out of the specifications of the invention. Thealloy powder CA-3 has an N content and a total content of “Ti+Zr+Nb+Ta”that are outside of the specifications of the invention. The alloypowder CA-4 has a total content of “Ti+Zr+Nb+Ta” that fail to satisfythe specifications of the invention. The alloy powder CA-5 has a Wcontent and a total content of “Ti+Zr+Nb+Ta” that are out of thespecifications of the invention.

Experimental 2

(Manufacturing of SLM Alloy Product Formed of IA-2 and SLM Alloy ProductFormed of CA-5)

AM articles (8 mm in diameter×10 mm in height) were formed of the alloypowders IA-2 and CA-5 with the particle size S prepared in Experimental1 by SLM (the selective laser melting step S2). The thickness of eachalloy powder bed h and the output power of the laser beam P were set at100 μm and 100 W, respectively. The local heat input P/S (unit:W×s/mm=J/mm) was controlled by varying the scanning speed (mm/s) of thelaser beam S. Controlling the local heat input corresponds tocontrolling the cooling rate.

Each AM article formed above was subjected to heat treatment at 1,150°C. with a holding duration of 4 hours (the solution heat treatment stepS3). Then, each solution heat-treated AM article was subjected to heattreatment at 900° C. with a holding duration of 4 hours (the aging heattreatment step S4) to manufacture an SLM alloy product formed of thepowder IA-2 or the powder CA-5.

(Manufacturing of LMD Alloy Product Formed of IA-2 and LMD Alloy ProductFormed of CA-5)

An AM article was formed of each of the alloy powders IA-2 and CA-5 withthe particle size L prepared in Experimental 1 by laser metal deposition(LMD). Each AM article was subjected to the solution heat treatment stepS3 and the aging heat treatment step S4 in a similar manner to the aboveto manufacture an LMD alloy product formed of the powder IA-2 or anotherLMD alloy product formed of the powder CA-5. The LMD process wasperformed with the output power of the laser beam P set at 800 W and thescanning speed of the laser beam S set at 15 mm/s.

Meanwhile, LMD is an AM process to form a deposit in which alloy powderis fed as a laser beam is irradiated. Generally, the local heat input ofLMD is larger than that of SLM. In other words, the cooling rate of LMDis slower than that of SLM.

(Manufacturing of Casting Alloy Product Formed of IA-2 and Casting AlloyProduct of CA-5)

A cast article (8 mm in diameter×10 mm in height) was formed of each ofthe alloy powder IA-2 and the alloy powder CA-5 with the particle size Lprepared in Experimental 1 by precision casting. Each cast article wassubjected to the solution heat treatment step S3 and the aging heattreatment step S4 in a similar manner to the above to manufacture a castalloy product formed of the alloy powder IA-2 or another cast alloyproduct formed of the alloy powder CA-5.

(Microstructure Observation and Mechanical Properties Testing)

Test pieces for microstructure observation and mechanical propertiestesting were taken from the AM articles, the cast articles, and theproducts manufactured above and subjected to microstructure observationand mechanical properties testing.

The microstructure observation was performed by SEM. Also, the obtainedSEM images were subjected to image analysis using an image processingprogram (ImageJ, a public domain program developed at the NationalInstitutes of Health, or NIH) to measure the average size of segregationcells, the average spacing of microsegregation, and the averageintergrain distance between carbide phase grains.

For the mechanical properties testing, a tensile test was performed atroom temperature (approximately 23° C.) to measure the 0.2% proofstress.

FIG. 5 is an SEM image showing an exemplary microstructure of a Co basedalloy AM article formed by LMD. FIG. 6 is an SEM image showing anexemplary microstructure of a Co based alloy cast article formed byprecision casting. Also, FIG. 2 above is an SEM image showing anexemplary microstructure of a Co based alloy AM article formed by SLM.The samples shown in FIGS. 2, 5, and 6 are formed of the alloy powderIA-2.

As mentioned before, the AM article formed by SLM (see FIG. 2) is apolycrystalline body that has segregation cells of about 1 μm in sizeformed in its crystal grains. In contrast, the AM article formed by LMD(see FIG. 5) is a polycrystalline body that has segregation cells ofabout 5 to 20 μm in size constituting each of its crystal grain. In thecast article formed by precision casting (see FIG. 6), microsegregationis observed at its dendrite boundaries, and the spacing of themicrosegregation is about 100 to 300 μm.

The microstructure observation of the products formed by subjecting theAM articles and the cast articles to the solution and aging treatmentsand the measurement of the average intergrain distance of carbide phasegrains revealed that the average intergrain distance of each productroughly matched the average size of segregation cells or the averagemicrosegregation spacing (the microstructures are not shown). Also, itwas found that when the average size of segregation cells is extremelysmall (e.g. about 0.1 μm or smaller), the solution treatment and theaging treatment caused adjacent carbide phase grains to combine to formlarger grains (as a result, the average intergrain distance betweencarbide phase grains widened).

Since the products obtained had largely different microstructures, arelationship between the average size of segregation cells and amechanical property was investigated. FIG. 7 is a graph showing anexemplary relationship between the average size of segregation cells ofthe Co based alloy AM articles and the 0.2% proof stress of the Co basedalloy products. FIG. 7 also shows the data of the cast articles and thecast alloy products for comparison. Regarding the cast articles, theaverage microsegregation spacing was used in place of the average sizeof segregation cells.

As shown in FIG. 7, the Co based alloy products formed of the powderCA-5 exhibit an almost constant 0.2% proof stress without being affectedby the average size of segregation cells. By contrast, the Co basedalloy products formed of the powder IA-2 largely vary in the 0.2% proofstress according to the average size of segregation cells.

The powder CA-5 has a total content of “Ti+Zr+Nb+Ta” that is too small(i.e. it contains almost none of them). Therefore, the products formedof the powder CA-5 have microstructures in which no reinforcing carbidephase is precipitated, but Cr carbide grains are precipitated. Thisconfirms that Cr carbide grains are not very effective as precipitationstrengthening grains. In contrast, the products formed of the powderIA-2 have microstructures in which reinforcing carbide phase grains areprecipitated. This is believed to have caused the substantial variationin 0.2% proof stress depending on the average segregation cell size(which determines the average intergrain distance between carbide phasegrains).

Also, considering the prescribed properties of turbine high temperaturecomponents toward which the invention is directed, a 0.2% proof stressof 500 MPa or more is required. So, when any 0.2% proof stress of 500MPa or more is judged as “Passed”, and any 0.2% proof stress of lessthan 500 MPa is judged as “Failed”, it has been confirmed thatmechanical properties that qualify as “Passed” are obtained with theaverage segregation cell size (i.e. the average intergrain distancebetween carbide phase grains) ranging from 0.15 to 1.5 μm. In otherwords, conventional carbide phase-precipitated Co based alloy materialshave failed to exhibit sufficient mechanical properties partly becausethe average intergrain distance between reinforcing carbide phase grainshas not been successfully controlled to fall within a desired range.

Experimental 3

(Manufacturing of SLM Alloy Products IP-1-1 to IP-5-1 Formed of IA-1 toIA-5 and SLM Alloy Products CP-1-1 to CP-5-1 Formed of CA-1 to CA-5)

An AM article (8 mm in diameter×10 mm in height) was formed of each ofthe alloy powders IA-1 to IA-5 and CA-1 to CA-5 of the particle size Sprepared in Experimental 1 by SLM (the selective laser melting step S2).Based on the results obtained in Experimental 2, the SLM process wasperformed with the average intergrain distance between carbide phasegrains controlled to 0.15 to 1.5 μm.

Each AM article formed above was subjected to heat treatment at 1,150°C. with a holding duration of 4 hours (the solution heat treatment stepS3). Then, the solution heat-treated AM articles were each subjected toheat treatment within a range from 750 to 1,000° C. with a holdingduration of 0.5 to 10 hours (the aging heat treatment step S4) tomanufacture SLM alloy products IP-1-1 to IP-5-1 formed of the powdersIA-1 to IA-5 and SLM alloy products CP-1-1 to CP-5-1 formed of thepowders CA-1 to CA-5.

(Microstructure Observation and Mechanical Properties Testing)

Test pieces for microstructure observation and mechanical propertiestesting were taken from the SLM alloy products IP-1-1 to IP-5-1 andCP-1-1 to CP-5-1 and subjected to microstructure observation andmechanical properties testing.

The microstructure observation was performed through image analysis ofSEM images in a similar manner to Experimental 2 to measure the averagematrix crystal grain size and the average intergrain distance betweencarbide phase grains.

As to the mechanical properties testing, a creep test was conducted at900° C. under a stress of 98 MPa to measure the creep rupture time.Based on the prescribed properties of turbine high temperaturecomponents toward which the invention is directed, any creep rupturetime of 1,100 hours or more was judged as “Passed”, and any creeprupture time of less than 1,100 hours was judged as “Failed”. When thecreep properties are judged as “Passed”, the temperature at which thecreep rupture time reaches 100,000 hours under a stress of 58 MPa is875° C. or higher. Such creep properties can be deemed as comparable tothose of Ni based alloy materials.

The results of Experimental 3 are shown in Table 2.

TABLE 2 Measurement and Testing Results of SLM Alloy Products IP-1-1 toIP-5-1 and CP-1-1 to CP-5-1. Average Intergrain Distance Between AlloyAlloy Average Matrix Reinforcing Carbide Product Powder Crystal GrainSize Phase Grains No. No. (μm) (μm) Creep Test IP-1-1 IA-1 75 0.6 PassedIP-2-1 IA-2 90 0.3 Passed IP-3-1 IA-3 60 0.8 Passed IP-4-1 IA-4 95 1.1Passed IP-5-1 IA-5 85 1.5 Passed CP-1-1 CA-1 90 Cr Carbide Phase FailedPrecipitation CP-2-1 CA-2 65 2.5 Failed CP-3-1 CA-3 75 Cr Carbide PhaseFailed Precipitation CP-4-1 CA-4 95 No Reinforcing Failed Carbide PhaseCP-5-1 CA-5 70 No Reinforcing Failed Carbide Phase

As shown in Table 2, the SLM alloy products IP-1-1 to IP-5-1 all passedthe creep testing. This is not only because the average matrix crystalgrain size is within the appropriate range but also because the averageintergrain distance between reinforcing carbide phase grains (MC typecarbide phase grains of Ti, Zr, Nb, and/or Ta) is sufficiently small(i.e. the reinforcing carbide phase grains are finely and dispersedlyprecipitated).

On the other hand, the SLM alloy products CP-1-1 to CP-5-1 all failedthe creep testing, although the average matrix crystal grain size waswithin the appropriate range. To determine possible causes behind this,the SLM alloy products CP-1-1 to CP-5-1 were examined individually.Regarding CP-1-1, the excessive contents of C and Cr have resulted indominant precipitation of Cr carbide grains. Regarding CP-2-1, theexcessive content of C and the excessive total content of “Ti+Zr+Nb+Ta”have resulted in coarsening of the reinforcing carbide phase grains andan increased average intergrain distance. Regarding CP-3-1, theexcessive content of C and the insufficient total content of“Ti+Zr+Nb+Ta” have resulted in dominant precipitation of the Cr carbidegrains. These results confirm that Cr carbide grains are not veryeffective as precipitation strengthening grains. Regarding CP-4-1 andCP-5-1, the insufficient total content of “Ti+Zr+Nb+Ta” (almost none)has resulted in no precipitation of the reinforcing carbide phaseitself.

Based on the results of Experimental 3, it has been confirmed that IA-1to IA-5, which have the chemical compositions specified in theinvention, are desirable as starting materials for SLM alloy products.It has also been confirmed that creep properties can be improved bycontrolling the average intergrain distance between reinforcing carbidephase grains to 0.15 to 1.5 μm.

Experiment 4

(Manufacturing of SLM Alloy Products IP-1-2 to IP-1-7 and IP-2-2 toIP-2-7)

AM articles (8 mm in diameter×10 mm in height) were formed of the alloypowders IA-1 and IA-2 of the particle size S prepared in Experimental 1by SLM (the selective laser melting step S2). Based on the resultsobtained in Experimental 2, the SLM process was performed with theaverage intergrain distance between carbide phase grains controlled to0.15 to 1.5 μm.

Each AM article formed above was subjected to the solution treatment andthe aging treatment. By varying the temperature and the holding durationof the solution treatment within ranges from 1,000 to 1,300° C. and from0.5 to 10 hours, respectively, SLM alloy products IP-1-2 to IP-1-7 andIP-2-2 to IP-2-7 varying in average matrix crystal grain size weremanufactured. The aging treatment conditions were set to be the same asExperimental 3.

(Microstructure Observation and Mechanical Properties Testing)

Test pieces for microstructure observation and mechanical propertiestesting were taken from the SLM alloy products IP-1-2 to IP-1-7 andIP-2-2 to IP-2-7 and subjected to microstructure observation andmechanical properties testing.

The microstructure observation was performed through image analysis ofSEM images in a similar manner to Experimental 2 to measure the averagematrix crystal grain size. Also, as mechanical properties testing, acreep test was conducted in a similar manner to Experimental 2, and eachproduct was judged as “Passed” or “Failed” based on the same criteria asExperimental 2. The results of Experiment 4 are shown in Table 3.

TABLE 3 Measurement and Testing Results of SLM Alloy Products IP-1-2 toIP-1-7 and IP-2-2 to IP-2-7. Alloy Alloy Average Matrix Product PowderCrystal Grain Size No. No. (μm) Creep Test IP-1-2 IA-1 11 Failed IP-1-315 Failed IP-1-4 43 Passed IP-1-5 74 Passed IP-1-6 151 Failed IP-1-7 180Failed IP-2-2 IA-2 15 Failed IP-2-3 32 Passed IP-2-4 56 Passed IP-2-5 74Passed IP-2-6 141 Passed IP-2-7 200 Failed

As shown in Table 3, it has been confirmed that the average matrixcrystal grain size is preferably 20 to 145 μm. Also, based on theresults of Experimental 4, it has been confirmed that the solutiontreatment is preferably performed within a temperature range of 1,100 to1,200° C. with a holding duration of 0.5 to 10 hours.

Experiment 5

(Examination of SLM Conditions in Selective Laser Melting Step)

AM articles (8 mm in diameter×10 mm in height) were formed of the alloypowders IA-4 of the particle size S prepared in Experimental 1 by SLM(the selective laser melting step S2). The output power of the laserbeam P was set at 85 W, and the local heat input P/S (unit: W×s/mm=J/mm)was controlled by varying the thickness of the alloy powder bed h andthe scanning speed (mm/s) of the laser beam S.

The AM articles formed above were each subjected to microstructureobservation to measure the average segregation cell size. As withExperimental 2, the microstructure observation was conducted by SEM andthe measurement was conducted using an image processing program(ImageJ).

FIG. 8 shows exemplary SLM conditions to obtain a Co based alloy AMarticle according to the embodiment of the invention, indicating arelationship between the thickness of the alloy powder bed and the localheat input. In FIG. 8, “o” signifies the AM articles observed to have anaverage segregation cell size within a range of 0.15 to 1.5 μm andjudged as “Passed”, and “x” signifies the other AM articles, judged as“Failed”.

Based on the results of Experimental 5, it has been confirmed that inthe selective laser melting step S2, the SLM process is preferablyperformed while controlling the thickness of the alloy powder bed h(unit: μm), the output power of the laser beam P (unit: W), and thescanning speed of the laser beam S (unit: mm/s) such that they satisfythe following formulas: “15<h<150” and “67×(P/S)−3.5<h<2222×(P/S)+13”.In other words, the hatched region is the Passed region.

The above-described embodiments and Examples have been specificallygiven in order to help with understanding on the present invention, butthe invention is not limited to the described embodiments and Examples.For example, a part of an embodiment may be replaced by known art, oradded with known art. That is, a part of an embodiment of the inventionmay be combined with known art and modified based on known art, as faras no departing from a technical concept of the invention.

What is claimed is:
 1. A cobalt based alloy article, having a chemicalcomposition comprising: 0.08 to 0.25 mass % of carbon; 0.1 mass % orless of boron; 10 to 30 mass % of chromium; 5 mass % or less of iron, 30mass % or less of nickel, the total amount of the iron and the nickelbeing 30 mass % or less; tungsten and/or molybdenum, the total amount ofthe tungsten and the molybdenum being 5 to 12 mass %; titanium,zirconium, niobium and tantalum, the total amount of the titanium, thezirconium, the niobium and the tantalum being 0.69 to 2 mass %; 0.5 mass% or less of silicon; 0.5 mass % or less of manganese; 0.003 to 0.04mass % of nitrogen; and the balance being cobalt and impurities, andwherein in crystal grains of the cobalt based alloy article, there aresegregation cells having an average size of 0.15-1.5 microns.
 2. Thecobalt based alloy article according to claim 1, wherein the chemicalcomposition of the cobalt based alloy comprises 0.01 to 1 mass % of thetitanium, 0.05 to 1.5 mass % of the zirconium, 0.02 to 1 mass % of theniobium, and 0.05 to 1.5 mass % of the tantalum.
 3. The cobalt basedalloy article according to claim 1, wherein the chemical composition ofthe cobalt based alloy comprises, as the impurities, 0.5 mass % or lessof aluminum and 0.04 mass % or less of oxygen.
 4. A cobalt based alloyarticle, having a chemical composition comprising: 0.08 to 0.25 mass %of carbon; 0.1 mass % or less of boron; 10 to 24.5 mass % of chromium; 5mass % or less of iron, 30 mass % or less of nickel, the total amount ofthe iron and the nickel being 30 mass % or less; tungsten and/ormolybdenum, the total amount of the tungsten and the molybdenum being 5to 12 mass %; titanium, zirconium, niobium and tantalum, the totalamount of the titanium, the zirconium, the niobium and the tantalumbeing 0.5 to 2 mass %; 0.5 mass % or less of silicon; 0.5 mass % or lessof manganese; 0.003 to 0.04 mass % of nitrogen; and the balance beingcobalt and impurities, and wherein in crystal grains of the cobalt basedalloy article, there are segregation cells having an average size of0.15-1.5 microns.